Introduction

A new class of entropic alloys (also referred to as nonlinear alloys) was first proposed by Yeh et al. in 1995, and alloys with outstanding properties were prepared in 2004 (Ref 1). HEAs may be defined as containing multiple principle elements, and the atomic percentages of each element range between 5 and 35% (Ref 2). Due to the benefits of the high entropy of mixing, HEAs have the ability to resist the effects of mixing enthalpy and restrain the precipitation of ordered intermetallic compounds. Therefore, when cooling from the molten state, HEAs are intended to form simple, disordered, face-centered cubic (FCC), body-centered cubic (BCC), or hexagonal close-packed (HCP) solid solutions rather than ordered phases (Ref 3,4,5,6,7,8). And, according to the previous study results, the HEAs can possess comprehensive properties which may not be available in other traditional alloys (Ref 9,10,11,12,13,14,15,16,17,18,19,20,21,22,23,24,25,26). Therefore, since they provide a new alloy designing concept, the HEAs may be marked as a milestone in the development of structural and functional metallic materials. In particular, due the aforementioned positive properties, the HEAs have greatly expanded the current vast applicable alloy systems.

The low-density HEAs which have been investigated in recent years are considered to be important lightweight materials with high strength (Ref 27, 28). Youssef et al. studied low-density nanostructured HEAs (grain size: approximately 12 nm) Li20Al20Mg10Sc20Ti30 using a mechanical alloying technique (Ref 29). The alloys exhibited a single FCC phase structure, with a high hardness (5.9 GPa). It was observed that after annealing at 500 °C, the crystal structures transformed to a HCP phase structures, and the hardness decreased to 4.9 GPa. In order to meet the demands of the aerospace industry, Senkov et al. (Ref 30, 31) designed low-density refractory Cr-Nb-Ti-V-Zr alloy systems using vacuum arc melting, followed by hot isostatic pressing (HIP) and homogenization annealing, which exhibited densities in the range of 6.34 g/cm−3 (NbTiV2Cr) to 6.67 g/cm−3 (CrNbTiZr). They found that the CrNiTiVZr alloy had more attractive properties than that the In718 and Haynes 230 alloys. However, the density was found to be significantly high, and further weight reduction was required for lightweight applications. An AlTiVCr equiatomic alloy with a single-phase structure was prepared by Qiu et al. (Ref 32). The density of the alloy was 5.06 g/cm3. In another related study, Li et al. examined Mgx(AlCuMnZn)100-x HEAs using induction melting and casting in copper molds (Ref 33, 34). The alloys exhibited mainly HCP phases and Al-Mn icosahedral quasicrystalline phases, with densities ranging from 2.20 to 4.29 g/cm−3 depending on the percentage of Mg in the composition. Those alloys showed the compression strengths of 400 to 500 MPa and the ductilities of 3-5% at room temperature. Beak et al. reported the microstructures and compressive properties of Al70Mg10Si10Cu5Zn5 alloys at room and elevated temperatures (Ref 35,36,37). An ultrasonic melt treatment (UST) was used to improve the ultimate compressive strength of as-cast alloys from 573 to 681 MPa, since the USTs could effectively decrease both the grain sizes and precipitated phase sizes. Then, the precipitation behaviors of the Al-6Mg-9Si-10Cu-10Zn-3Ni alloys were investigated, along with their impacts on the compressive strength. The mechanical properties were found to be greatly improved, causing fine clusters and precipitates to be formed at below 70 °C. The Al-6Mg-9Si-10Cu-10Zn-3Ni alloys showed excellent mechanical properties (> 700 MPa) when compared to some of the Al alloys (A356 and A390). In another related study conducted by Yang et al. (Ref 38), heavy elements were substituted for the Al, Li, and Mg light elements in Al-Li-Mg systems. It was found that the densities varied from 2.84 g/cm3 (Al80Li5Mg5Zn5Sn5) to 3.88 g/cm3 (AlLiMgZnSn). However, the resulting microstructures were multi-phased, with numerous intermetallic phases observed.

In order to develop advanced lightweight alloys, a series of low-density multicomponent alloys based on the Al-Mg system were investigated according to a high-entropy strategy (Ref 2, 39,40,41,42,43,44). In accordance with the results of the previous studies (Ref 38), the content levels of zinc, copper, and silicon were designed to a 10% mole fraction. In this study, from the perspective of a ternary phase diagram, a straight line could be drawn which joined two points, Al62Mg38 and ZnCuSi. The schematic is shown in Fig. 1. The point of the Al62Mg38 was decided due to its eutectic point of Al-Mg system alloys. The eutectic point displayed the best fluidity and casting properties. The aforementioned line intersected the line of 10 atomic percent (at.%) content of zinc, copper, and silicon at the A1 (Al58.5Mg31.5Zn4.5Cu4.5Si1) point. Next, the other two points, A2 (Al63Mg27Zn4.5Cu4.5Si1) and A3 (Al66.7Mg23.3Zn4.5Cu4.5Si1), were chosen in the line of 10 at.% content of zinc, copper, and silicon. After examining the comprehensive mechanical properties of the three points, it was found that the A3 point displayed better properties than the other two points. Therefore, points Al and A3 were connected, and the content of B1 (Al80Mg14Zn2.7Cu2.7Si0.6), B2 (Al85Mg10.5Zn2.025Cu2.025Si0.45), and B3 (Al90Mg7Zn1.35Cu1.35Si0.3) was selected. The microstructure, phase composition, and mechanical behaviors were described herein.

Fig. 1
figure 1

Design schematic of Al-Mg-based multicomponent alloy

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Experimental Procedures

In this study, the Al-Mg-based medium entropic alloys were produced by the vacuum induction melting of the simple mixtures of the corresponding elements under a high-purity argon atmosphere in a graphite crucible. All of the raw materials were in bulk form, with a purity of 99.9 weight percent (wt.%). Prior to placing in the crucible, all of the elements were polished using a grinding machine in order to remove the oxide film. Then, to achieve a homogeneous distribution of the elements in the alloys, each alloy was re-melted several times in order to guarantee a total of 30 min in a liquid state. The molten alloys were then cast into a Φ75 mm × 10 mm steel mold.

All of this study’s tests were carried out in the as-cast condition. The phase formation was determined by x-ray diffraction (XRD) using a SmartLab diffractometer with Cu-α radiation. Then, specimens with dimensions of 8 × 8 × 10 mm3 were ground, polished, and etched by a Keller reagent for subsequent microstructure observations. The microstructures were characterized by a ZEISS SUPRA 55 scanning electron microscope (SEM) which was installed with an energy-dispersive spectrometer (EDS). Compression tests with a strain rate of 1 × 10−3 s−1 at room temperature were carried out on cylindrical samples (3 mm in diameter and 6 mm in length), which had been cut from the caste alloys using an electric discharged machine. For the compression tests, the cylindrical samples were carefully polished, and at least three specimens were tested for each alloy in order to obtain the statistical distributions of the tensile properties. The densities of the alloys were measured using an Archimedes method.

Results

Densities of the Alloys

The densities (ρ) of the targeted alloys were measured using an Archimedes method, as detailed in Table 1.

Table 1 Chemical compositions (at.%) and densities (ρ) of alloys in this work
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A rule of mixture assumption of a solid solution was used for examining the targeted alloys, and the estimated theoretical densities (ρmix) are listed in Table 1 (Ref 45,46,47):

$$ \rho_{\text{mix}} = \frac{{\sum\nolimits_{i = 1}^{n} {c_{i} A_{i} } }}{{\sum\nolimits_{i = 1}^{n} {c_{i} A_{i} /p_{i} } }} $$
(1)

where ci, Ai, and ρi are the weight fraction, atomic weight, and density of the ith corresponding element, respectively, and n denotes the total number of elements. It was found that when comparing the measured density and corresponding theoretical mixed densities, significant deviations were observed, which suggested that ordered intermetallic compounds were present.

Phase Formations and Microstructures

The results of this study’s x-ray diffraction and corresponding backscattered electron images are shown in Fig. 2 and 3, respectively. All of the major phases are identified in Fig. 2. Meanwhile, the phase compositions of the alloys were found to be complex, not only in regard to the solid solution phases, but also the intermetallic phases. Furthermore, the scanning electron micrographs were also complex, with at least two constituents observed in each alloy. Moreover, remarkable differences were observed between the fabricated alloys. The compositions of the phases measured by the EDS analysis are listed in Table 2.

Fig. 2
figure 2

XRD patterns of as-cast lightweight Al-Mg-based multicomponent alloys

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Fig. 3
figure 3

SEM backscattered electron images of as-cast lightweight Al-Mg-based multicomponent alloy

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Table 2 Chemical compositions (at.%) of phases identified in the microstructures of Al55.8Mg34.2Zn4.5Cu4.5Si1, Al63Mg27Zn4.5Cu4.5Si1,Al66.7Mg23.3Zn4.5Cu4.5Si1, Al80Mg14Zn2.7Cu2.7Si0.6, Al85Mg10.5Zn2.025Cu2.025Si0.45, and Al90Mg7.1Zn1.35Cu1.35Si0.3
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Figure 2(a) details the XRD pattern of the A1 alloy. When comparing the alloy with traditional HEAs with BCC and/or FCC phases, it was not surprising to find them to be more complex (Ref 48,49,50,51). The ordered Mg32(AlZn)49 and Mg2Si phases, along with a solid solution phase, were observed in the diffraction graph of the alloy. In this study, the intensity of the diffraction peak of the ordered intermetallic phase was considered as the strongest and was obviously higher than the second and the third peaks of the FCC solid solution phases, which indicated that the intermetallic phase was the dominant phase. The SEM images of the A1 alloy are shown in Fig. 3(a), in which five constituents (marked as A, B, C, D, and E, respectively) with different contrasts are evident. The phase constituents of the various regions were identified by EDS analysis, and the results are listed in Table 2. Then, combined with the EDS analysis results, the phase constituents of each region in the A1 alloy were successfully identified. It was believed that Region A was most probably the Mg32(AlZn)49 phase, and Region B was a matrix of the α-Al based solid solution. The black rhombic (marked as C in the figure) denoted the Mg2Si phase. Region D was the eutectic phase, and Region E contained the segregation elements of Cu and Zn.

The XRD pattern of the A2 alloy is shown in Fig. 2(b). The A2 alloy was similar to the A1 alloy and composed of the FCC phase, Mg32(AlZn)49 and Mg2Si. However, the intensities of ordered phases were weaker than those detailed in Fig. 2(a), which suggested that the volume fraction of the ordered phases had been decreased. The microstructure of the A2 alloy is presented in Fig. 3(b), in which four different constituents are evident and marked as A, B, C, and D. In this study, in accordance with the microstructure and EDS results listed in Table 2, it was indicated that Region A was probably the Mg2Si; Region B was a α-Al based solid solution; and Region D was a typical eutectic structure. When compared with the A1 alloy, the eutectic phase in the A2 alloy displayed a higher volume, which may have been the result of the constituent of the A2 alloy having a close proximity to the eutectic point.

When the content of the aluminum reached 66.7 at.%, there are no changes observed in the types of phases in the A3 alloy. However, there were decreases observed in the peaks intensities of the ordered intermetallic phases, while the FCC phase became notable (Fig. 2c). There were four distinct alloy constituents found, marked as A, B, C, and D in Fig. 3(c). In accordance with the EDS results shown in Table 2, it was concluded that Region A corresponded to a α-Al solid solution phase; Region B was Mg2Si; and Regions C and D were similar to Regions C and D in the A2 alloy, respectively.

The XRD patterns of the Al-Mg-rich alloys, B1, B2, and B3 are shown in Fig. 2(d) to (f), respectively. There were two phases [α-Al and Mg32(AlZn)49] recognized in the XRD patterns of the Al-Mg-rich alloys. It was found that with the increases in aluminum content, the intensities of the peaks of the α-Al became stronger, while those of the Mg32(AlZn)49 became weaker. Eventually, only the α-Al phase could be identified in the B3 alloys containing 90 at.% aluminum. Those alloys were determined to be composed of the FCC phase as the major phase, and the ordered phase of the Mg32(AlZn)49 or unknown phases were considered to be minor phases. The SEM images of the aforementioned alloys are presented in Fig. 3(d) to (f), respectively. The Al-Mg-rich alloys displayed thick dendritic structures. It was observed that the higher the aluminum content was, the thicker the dendrite would be, and the matrix was divided by the eutectic structure. When combined with the EDS analysis results listed in Table 2, the eutectic regions were determined to be rich in Mg, Cu, and Zn, which may have been Mg32(AlZn)49 or other unknown phases. However, the B regions in Fig. 3(d) to (f) were found to be Cu enriched. This trend suggested that, due to a higher melting temperature, the element Cu had segregated during the solidification process.

Mechanical Behavior

Figure 4 details the compressive stress–strain curves for the as-cast alloys at room temperature. The engineering properties of these alloys are listed in Table 3. During this study’s compression testing at room temperature, the A1, A2, and A3 alloys failed without obvious yielding. It was found that they had fractured by shattering into many pieces after approaching the maximum strengths of 577 MPa, 676 MPa, and 660 MPa at the engineering strains of 3.2, 1.9, and 4.2%, respectively. The test results revealed that these alloys exhibited high compressive strength (exceeding 550 MPa), but very low plasticity. The room temperature compression behaviors of the B1, B2, and B3 alloys were found to be different from those described above. These alloys displayed plastic yielding at σ0.2 = 203, 255, and 198 MPa, respectively. After that, a continuous increase in their strengths occurred with increases in the compression strain. For example, the strengths of the B2 and B3 alloys had increased to σ20 = 721 MPa and 669 MPa, respectively, after compression by 20% at room temperature. It was observed that with the increases in the aluminum content, the engineering strain had greatly improved from 1.9 to 32.7%. The B1 alloy fracture strength (σf) had decreased to 498 MPa. However, that of the B2 and B3 alloys had approached or exceeded 800 MPa and displayed excellent room temperature compressive strengths. In other words, the plasticity of these alloys had been improved by increasing the aluminum content and the volume fraction of the more ductile α-Al phase.

Fig. 4
figure 4

Compressive stress–strain curves at room temperature. The strain rate is 10−3 s−1

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Table 3 Compressive mechanical properties of Al–Mg-Zn-Cu-Si alloys
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Discussion

From the viewpoint of the thermodynamic concept, a system will attempt to minimize its Gibbs free energy (ΔGmix) in order to achieve a metastable or stable equilibrium state, which is known to be related to the enthalpy (∆Hmix) and entropy (∆Smix). The Gibbs free energy formula is shown in Eq 2:

$$ \Delta G_{\text{mix }} = \Delta H_{\text{mix }} - T\Delta S_{\text{mix }} $$
(2)

It can be seen that the Gibbs free energy (ΔGmix) is determined by the enthalpy of the mixing (ΔHmix) and the entropy of the mixing (ΔSmix) at a given temperature (T) for an equilibrium state. It is the competition result between the ΔHmix and TΔSmix, which affects the phase selection in the HEAs.

Furthermore, the mixing enthalpy and entropy are two important parameters in the study of HEAs and multicomponent alloys. Therefore, based on the thermodynamic and Hume–Rothery rules for solid solution formations, the phase formations in HEAs can be characterized by some specifically derived parameters, such as the ΔHmix (enthalpy of mixing); δ (atomic size difference); Ω (ratio of entropy to enthalpy); VEC (valance electron concentration); and Δχ (Pauling electronegativity difference). These criteria and related parameters are defined as follows:

$$ \Delta S_{\text{mix }} = - R\mathop \sum \limits_{i = 1}^{n} c_{i } { \ln }c_{i} $$
(3)

where ∆Smix is the mixing entropy; R is the gas constant; and ci denotes the atomic percentages of the ith element. The enthalpy of the mixing ∆Hmix can be determined as follows:

$$ \Delta H_{\text{mix}} = \mathop \sum \limits_{i = 1,j = 1,i \ne j}^{n} \varOmega_{ij} c_{i} c_{j} $$
(4)

where Ωij= 4ΔHmix AB; ΔHmix AB is the mixing of the enthalpy of binary equiatomic AB alloys; and Ωij is the regular melt interaction parameter between the ith and jth elements. Then, the atomic size difference (δ) can be obtained by the following:

$$ \delta = \sqrt {\mathop \sum \limits_{i = 1}^{n} c_{i} \left( {1 - \frac{{r_{i} }}{{\bar{r}}}} \right)^{2} } $$
(5)

where n is the number of the involved elements in an alloy system; ri denotes the atomic radius; and \( \bar{r} = \mathop \sum \limits_{i = 1}^{n} c_{i} r_{i} \) is the average atomic radius. Zhang et al. proposed a parameter (Ω) to predict the phase formation (Ref 39, 52) as follows:

$$ \varOmega = \frac{{T_{m} \Delta S_{\text{mix}} }}{{\left| {\Delta H_{\text{mix}} } \right|}} $$
(6)

where \( T_{m} = \mathop \sum \limits_{i = 1}^{n} c_{i} \left( {T_{m} } \right)_{i} \) is the average melting temperature of the n-elements HEA; and (Tm)i is the melting temperature of the ith component. The Pauling electronegativity difference and valence electron concentration were as follows:

$$ \Delta \chi = \sqrt {\mathop \sum \limits_{i = 1}^{n} c_{i} \left( {\chi_{i} - \bar{\chi }} \right)^{2} } $$
(7)
$$ VEC = \mathop \sum \limits_{i = 1}^{n} c_{i } (VEC)_{i } $$
(8)

where \( \bar{\chi } = \mathop \sum \limits_{i = 1}^{n} c_{i} \chi_{i} \), χi is the Pauling electronegativity for the ith element; and (VEC)i is the VEC of the ith element.

Among all of the aforementioned criteria, the atomic size difference (δ) plays a crucial role in phase formations. A large value of δ will indicate low element diffusion and phase formation rates, after which the nanoparticles or amorphous phases are separated out, and the stability of a solid solution becomes broken down. In contrast, small values of δ promote the formations of solid solutions. In fact, a δ ≤ 6.6% has been considered as a criterion for the formation of solid solution phases in compositionally complex alloys. Zhang et al. summarized δ and ΔHmix in analyzed alloys and after plotting their relationships concluded that solid solutions tend to form in the region of − 22 kJ/mol ≤ ΔHmix ≤ 5 kJ/mol.

In this research study, in order to better understand phase formations in the low-density multicomponent alloys, the relationship between δ and the other four criterions was plotted, as detailed in Fig. 5. The corresponding phase constituents were derived from the research conducted by Zhang and Yang (Ref 33) and the value ranges of the criteria were calculated based on Eqs 3 to 8. Yang et al. believed that along with small δ values, the near-zero values of the absolute ΔHmix, large values of Ω (≥ 1.1), and small Δχ (≤ 0.175) effectively favored the stability of solid solutions, rather than those of intermetallic compounds. Therefore, δ-ΔHmix, δ-Ω, and δ-Δχ graphs could be used to predict the solid solution formations for HEAs which were based on 3d and/or 4d transition metal (TM) elements. Meanwhile, the VEC rule provided a convenient way to design the FCC or BCC HEAs containing mainly 3d and/or 4d TM elements. It was also pointed out that the configurational entropy was not enough to form a majority of the solid solution phase in low-density alloys containing a mass of Al, Mg, and other light elements. The critical values of the parameters were required to be modified in order to account for the apparent light elements, smaller values of δ (δ ≤ 4.5%), and greater ΔHmix (− 1 kJ/mol ≤ ΔHmix ≤ 5 kJ/mol), and Ω (Ω > 10). Meanwhile, the Δχ was determined to be the most predictable parameter across the extensive HEAs. Overall, the intermetallic compounds were determined to be stabilized when Δχ ≥ 0.175.

Fig. 5
figure 5

Phase constituent prediction maps (SS: solid solution; IC: intermetallic compound)

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As listed in Table 4, the values of the δ, ΔHmix, Ω, VEC, and Δχ were calculated in this study. Then, the relative locations of those alloys in the δ-ΔHmix, δ-Ω, δ-Δχ, and δ-VEC maps were plotted, as detailed in Fig. 5. It was clear these six lightweight alloys were inside the range of “traditional” solid solutions (containing mainly 3d and/or 4d TM elements), which confirmed that they were not in good agreement with the rules for low-density multicomponent alloys. However, the B1, B2, and B3 alloys mainly contained the aluminum element, and almost their α-Al phases are illustrated in Fig. 3(d) to (f). In particular, the B2, and B3 alloys displayed good mechanical behavior fitting with the criteria of δ and ΔHmix. Fortunately, the parameter Δχ was found to seemingly distinguish the solid solution formation behavior most uniformly. According to Yang’s view (Ref 33), the configurational entropy is not sufficient to form the majority of the solid solution phases in low-density alloys containing a mass of Al, Mg, and other main group elements. Therefore, it can be concluded that the δ, ΔHmix, and Δχ should potentially be regarded as the parameters for low-density multicomponent alloys, where δ ≤ 4.5%, − 1 kJ/mol ≤ ΔHmix ≤ 5 kJ/mol, and Δχ ≤ 0.175.

Table 4 Calculated values for δ, ΔHmix, ΔSmix, Tmix, Ω, VEC, and ∆χ of alloys in this work
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The mixing of enthalpy promotes the formation of the intermetallic compounds, while the mixing of entropy promotes the formation of solid solutions. The values of the mixing enthalpy of binary systems (ΔHmix AB) are listed in Table 5. The Mg-Si binary system is known to have the most negative mixing of enthalpy (− 26 kJ/mol), which means that the bonding force between these two elements is the strongest, and Mg2Si phase formation often occurs. Thereafter, with the Si element exhausted, the temperatures of molten alloys will be close to the solidification point of the Cu element. The EDS map observed that the Cu element appeared segregated at the grain boundary. This was likely due to the crystallization of a multicomponent eutectic, which included a copper-containing phase. It is interesting to note that the Cu element always occurs at a high content of the Mg element area, with the exception of the Mg2Si region. This phenomenon may be due to the fact that the mixing enthalpy between the Cu and Mg is higher than the others. The EDS analysis showed that the aluminum elements in the crystal and the grain boundary had the greatest degrees of composition fluctuations. For the main alloying elements, the most serious segregation was the Cu element, followed by the Zn element.

Table 5 Values of enthalpy of mixing ΔHmix AB (kJ mol−1), calculated by Miedema’s model for atomic pairs between constitutive elements of Al–Mg-Zn-Cu-Si alloys (in the upper right triangular region) and equilibrium phases for equiatomic binary systems contained in Al–Mg-Zn-Cu-Si alloys at room temperature (in the lower left triangular region)
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Table 6 details the volume fraction of intermetallic, eutectic, and solid solution phases of the alloys. The fracture strength of the Group A alloys was determined to be linearly related to the eutectic volume fraction, and the linear fitting equation was as follows:

$$ y \, = \, 334x \, + \, 441 $$
(9)

where x is the eutectic volume fraction and y denotes the fracture strength. With the increasing aluminum content, the element ratio reached the eutectic point, and a vast eutectic structure was formed, as shown in Fig. 3(a) and (b). The group of A alloys displayed a brittle fracture mode, and cracks were initiated from the boundary of the intermetallic. The eutectic structure was the mechanical mixture of the Mg32(AlZn)49 intermetallic and a solid solution and displayed a chrysanthemum distribution. The finely dispersed eutectic structure was found to be well distributed in all of the substrates, which was useful in preventing the extension of cracks and increasing strength. The A2 alloy exhibited the highest σf of the Group A alloys, as shown in Fig. 4 and Table 3. When the elemental composition reached the A3 point, the bulk solid solutions began to precipitate, as shown in Fig. 3(c). Then, with the increasing stress conditions, cracks were initiated and crossed the solid solution, which exhibited high plasticity. At this point, plastic deformation took place. Therefore, the A3 alloy had displayed a higher εp than the A1 and A2 alloys. The mechanical properties of the Group B alloys were observed to be different from those described above. The fracture strengths and solid solution fractions displayed a parabola curve relationship when the Al composition reached 80%. The parabola equation was as follows:

$$ y = - 165x^{2} + \, 2.61 \times 10^{4} x \, - 9479 $$
(10)

where x is the solid solution volume fraction and y represents the fracture strength. When the aluminum content reached 80%, the volume fraction of the solid solution was predominant, as shown in Fig. 3(d) to (f), and the Group B alloys displayed a ductile fracture mode. During the process of plastic deformation, a dislocation slip had occurred in the solid solution, and the slip process was blocked when the dislocation extended to the eutectic structure. The eutectic structure was able to impede the movement of the dislocation and also played a very important role in the strengthening of the alloys. It was observed that the more the volume fraction of the eutectic, the higher the σf. Also, the eutectic dispersion was found to be more diffuse in the substrate with higher σf. Therefore, there was determined to be a maximal influence of the eutectic fraction on the fraction strength. The composition of the B2 alloy was found to be closest to the maximum, and it also showed the highest σf of the Group B alloys.

Table 6 Volume fractions of intermetallic, eutectic, and solid solution phases in alloy
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Conclusions

In this research study, a series of lightweight Al-Mg-based entropic alloys were examined in-depth, which had been designed based on the concept of HEAs. It was found that all of the alloys exhibited low densities ranging from 2.64 to 2.75 g/cm3, which were below the density of standard titanium alloys (4.5 g/cm−3). The examined alloys consisted of various intermetallic phases and an α-Al solid solution phase in the as-cast state. When the aluminum concentration increased, the volume fraction of the intermetallic was observed to gradually decline, and the α-Al solid solution phase was dominant. The fabricated lightweight HEAs had high strength, with compressive strength exceeding 500 MPa at room temperature. The strength of the B2 alloy even surpassed 800 MPa. The Al-Mg-rich B1, B2, and B3 alloys displayed excellent plasticity, with high strain to failure values of 13.8, 24.8, and 32.7%, respectively. The critical values of the δ, ΔHmix, Ω, VEC, and Δχ were discussed in this study in order to explain the apparent differences in behavior of the lightweight element compositions. The modified δ, ΔHmix, and Δχ were regarded as the probable parameters for the lightweight multicomponent alloys, where δ ≤ 4.5%, − 1 kJ/mol ≤ ΔHmix≤ 5 kJ/mol, and Δχ ≤ 0.175.